Permanent magnet including multiple ferromagnetic phases and method for producing the magnet

ABSTRACT

An iron-based rare earth alloy magnet has a composition represented by the general formula: (Fe 1-m T m ) 100-x-y-z Q x R y M z , where T is at least one element selected from the group consisting of Co and Ni; Q is at least one element selected from the group consisting of B and C; R is at least one rare earth element substantially excluding La and Ce; and M is at least one metal element selected from the group consisting of Ti, Zr and Hf and always includes Ti. In this formula, the mole fractions x, y, z and m meet the inequalities of: 10 at %&lt;x≦20 at %; 6 at %≦y&lt;10 at %; 0.1 at %≦z≦12 at %; and 0≦m≦0.5, respectively.

BACKGROUND OF THE INVENTION

[0001] 1. Field of the Invention

[0002] The present invention generally relates to a method for producinga permanent magnet effectively applicable to motors and actuators ofvarious types, and more particularly relates to a method for producingan iron-based rare earth alloy magnet including multiple ferromagneticphases.

[0003] 2. Description of the Related Art

[0004] Recently, it has become more and more necessary to furtherenhance the performance of, and further reduce the size and weight of,consumer electronic appliances, office automation appliances and variousother types of electric equipment. For these purposes, a permanentmagnet for use in each of these appliances is required to maximize itsperformance to weight ratio when operated as a magnetic circuit. Forexample, a permanent magnet with a remanence B_(r) of 0.5 T or more isnow in high demand. Hard ferrite magnets have been used widely becausemagnets of this type are relatively inexpensive. However, the hardferrite magnets cannot show that high remanence B_(r) of 0.5 T or more.

[0005] An Sm—Co magnet, produced by a powder metallurgical process, iscurrently known as a typical permanent magnet with that high remanenceB_(r) of 0.5 T or more. Examples of other high-remanence magnets includeNd—Fe—B type magnets produced by a powder metallurgical or meltquenching process. An Nd—Fe—B type magnet of the former type isdisclosed in Japanese Laid-Open Publication No. 59-46008, for example,and an Nd—Fe—B type magnet of the latter type is disclosed in JapaneseLaid-Open Publication No. 60-9852, for instance.

[0006] However, the Sm—Co magnet is expensive, because Sm and Co areboth expensive materials.

[0007] As for the Nd—Fe—B type magnet on the other hand, the magnet ismainly composed of relatively inexpensive Fe (typically accounting for60 wt % to 70 wt % of the total quantity), and is much less expensivethan the Sm—Co magnet. Nevertheless, it is still expensive to producethe Nd—Fe—B type magnet. This is partly because huge equipment and agreat number of process steps are needed to separate and purify, or toobtain by reduction reaction, Nd, which usually accounts for 10 at % to15 at % of the total quantity. Also, a powder metallurgical processnormally requires a relatively large number of process steps by itsnature.

[0008] Compared to an Nd—Fe—B type sintered magnet formed by a powdermetallurgical process, an Nd—Fe—B type rapidly solidified magnet can beproduced at a lower process cost by a melt quenching process. However,to obtain a permanent magnet in bulk by a melt quenching process, abonded magnet should be formed by compounding a magnet powder, made froma rapidly solidified alloy, with a resin binder. Accordingly, the magnetpowder normally accounts for at most about 80 volume % of the moldedbonded magnet. Also, a rapidly solidified alloy, formed by a meltquenching process, is magnetically isotropic.

[0009] For these reasons, an Nd—Fe—B type rapidly solidified magnetproduced by a melt quenching process has a remanence B_(r) lower thanthat of a magnetically anisotropic Nd—Fe—B type sintered magnet producedby a powder metallurgical process.

[0010] As disclosed in Japanese Laid-Open Publication No. 1-7502, atechnique of adding at least one element selected from the groupconsisting of Zr, Nb, Mo, Hf, Ta and W and at least one more elementselected from the group consisting of Ti, V and Cr in combinationeffectively improves the magnetic properties of an Nd—Fe—B type rapidlysolidified magnet. When these elements are added, the magnet can haveits coercivity H_(cJ) and anticorrosiveness increased. However, the onlyknown effective technique of improving the remanence B_(r) is increasingthe density of a bonded magnet.

[0011] As for an Nd—Fe—B type magnet, an alternative magnet material wasproposed by R. Coehoorn et al., in J. de Phys, C8, 1998, pp. 669-670.The Coehoorn material has a composition including a rare earth elementat a relatively low mole fraction (i.e., around Nd_(3.8)Fe_(77.2)B₁₉,where the subscripts are indicated in atomic percentages) and an Fe₃Bprimary phase. This permanent magnet material is obtained by heating andcrystallizing an amorphous alloy that has been prepared by a meltquenching process. Also, the crystallized material has a metastablestructure in which soft magnetic Fe₃B and hard magnetic Nd₂Fe₁₄B phasescoexist and in which crystal grains of very small sizes (i.e., on theorder of several nanometers) are dispersed finely and uniformly as acomposite of these two crystalline phases. For that reason, a magnetmade from such a material is called a “nanocomposite magnet”. It wasreported that a nanocomposite magnet like this has a remanence B_(r) ofas high as 1 T or more. But the coercivity HE thereof is relatively low,i.e., in the range from 160 kA/m to 240 kA/m. Accordingly, thispermanent magnet is applicable only when the operating point of themagnet is 1 or more.

[0012] It has been proposed that various metal elements be added to thematerial alloy of a nanocomposite magnet to improve the magneticproperties thereof. See, for example, Japanese Laid-Open Publication No.3-261104, U.S. Pat. No. 4,836,868, Japanese Laid-Open Publication No.7-122412, PCT International Publication No. WO 00/03403 and W. C. Chanet. al., “The Effects of Refractory Metals on the Magnetic Properties ofα-Fe/R₂Fe₁₄B-type Nanocomposites”, IEEE Trans. Magn. No.5, INTER-MAG.99, Kyongiu, Korea, pp.3265-3267, 1999. However, none of these proposedtechniques can always obtain a sufficient “characteristic value percost”.

SUMMARY OF THE INVENTION

[0013] An object of the present invention is to provide a method forproducing an iron-based alloy permanent magnet, exhibiting excellentmagnetic properties including a high coercivity H_(cJ) of e.g., 480 kA/mor more and a high remanence B_(r) of e.g., 0.85 T or more, at a lowcost.

[0014] An iron-based rare earth alloy magnet according to the presentinvention has a composition represented by the general formula:(Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z), where T is at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare earth element substantially excluding La and Ce; and M is atleast one metal element selected from the group consisting of Ti, Zr andHf and always includes Ti. In this formula, the mole fractions x, y, zand m meet the inequalities of: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1at %≦z≦12 at %; and 0≦m≦0.5, respectively. The magnet has two or moreferromagnetic crystalline phases including hard and soft magneticphases. An average grain size of the hard magnetic phase is equal to orgreater than 10 nm and equal to or less than 200 nm, while that of thesoft magnetic phase is equal to or greater than 1 nm and equal to orless than 100 nm.

[0015] In one embodiment of the present invention, the mole fractions x,y and z preferably meet the inequalities of: 10 at %<x<17 at %; 8 at%≦y≦9.3 at %; and 0.5 at %≦z≦6 at %, respectively.

[0016] In another embodiment of the present invention, R₂Fe₁₄B phase,boride phase and α-Fe phase may coexist in the same metal structure.

[0017] Specifically, an average crystal grain size of the α-Fe andboride phases is preferably from 1 nm to 50 nm.

[0018] More specifically, the boride phase preferably includes aniron-based boride with ferromagnetic properties.

[0019] In this particular embodiment, the iron-based boride preferablyincludes Fe₃B and/or Fe₂₃B₆.

[0020] In still another embodiment, the mole fractions x and zpreferably meet the condition z/x≧0.1.

[0021] In yet another embodiment, the mole fraction y of the rare earthelement(s) R may be 9.5 at % or less.

[0022] Alternatively, the mole fraction y of the rare earth element(s) Rmay also be 9.0 at % or less.

[0023] In yet another embodiment, the magnet may have been shaped in athin strip with a thickness of 10 μm to 300 μm.

[0024] In yet another embodiment, the magnet may have been pulverizedinto powder particles.

[0025] Then, a mean particle size of the powder particles is preferablyfrom 30 μm to 250 μL m.

[0026] In yet another embodiment, the magnet may exhibit hard magneticproperties as represented by a coercivity H_(cJ) of 480 kA/m or more anda remanence B_(r) of 0.7 T or more.

[0027] In yet another embodiment, the magnet may also exhibit hardmagnetic properties as represented by a remanence B_(r) of 0.85 T ormore, a maximum energy product (BH)_(max) of 120 kJ/m³ or more and anintrinsic coercivity H_(cJ) of 480 kA/m or more.

[0028] A bonded magnet according to the present invention is formed bymolding a magnet powder, including the powder particles of the inventiveiron-based rare earth alloy magnet, with a resin binder.

[0029] A rapidly solidified alloy according to the present invention isa material for an iron-based rare earth alloy magnet. The alloy has acomposition represented by the general formula:(Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z), where T is at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare earth element substantially excluding La and Ce; and M is atleast one metal element selected from the group consisting of Ti, Zr andHf and always includes Ti. In this formula, the mole fractions x, y, zand m meet the inequalities of: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1at %≦z≦12 at %; and 0≦m≦0.5, respectively.

[0030] In one embodiment of the present invention, the rapidlysolidified alloy preferably has a structure, in which substantially noα-Fe phase is included but R₂Fe₁₄B compound and amorphous phases areincluded and in which the R₂Fe₁₄B phase accounts for a 60 volume percentor more.

[0031] Specifically, the mole fractions x, y and z preferably meet theinequalities of: 10 at %<x<17 at %; 8 at %≦y≦9.3 at %; and 0.5 at %≦z≦6at %, respectively. The R₂Fe₁₄B phase, accounting for 60 volume percentor more of the alloy, preferably has an average grain size of 50 nm orless.

[0032] Another rapidly solidified alloy according to the presentinvention is also a material for an iron-based rare earth alloy magnet.The solidified alloy is prepared by rapidly cooling a melt of a materialalloy comprising Fe, Q, R and Ti, where Q is at least one elementselected from the group consisting of B and C; and R is a rare earthelement. The solidified alloy has a structure in which an amorphousphase is included and in which heat treatment starts to grow a compoundcrystalline phase with an R₂Fe₁₄B crystalline structure before startingto grow an α-Fe crystalline phase.

[0033] An inventive method for producing an iron-based rare earth alloymagnet includes the steps of: preparing a melt of a material alloy thatincludes Fe, Q, R and Ti, where Q is at least one element selected fromthe group consisting of B and C, and R is a rare earth element; coolingthe melt to make a solidified alloy including an amorphous phase; andheating the solidified alloy to start growing a compound crystallinephase with an R₂Fe₁₄B crystalline structure and then an α-Fe crystallinephase.

[0034] In one embodiment of the present invention, the melt ispreferably cooled by a strip casting process.

[0035] Another inventive method for producing an iron-based rare earthalloy magnet includes the step of preparing a melt of a material alloy.The material alloy has a composition represented by the general formula:(Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z), where T is at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare earth element substantially excluding La and Ce; and M is atleast one metal element selected from the group consisting of Ti, Zr andHf and always includes Ti. The mole fractions x, y, z and m meet theinequalities of: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1 at %≦z≦12 at%; and 0≦m≦0.5, respectively. The method further includes the steps of:rapidly cooling the melt to make a rapidly solidified alloy in which anR₂Fe₁₄B crystalline phase and an amorphous phase coexist; andcrystallizing the rapidly solidified alloy to form a structure in whichtwo or more ferromagnetic crystalline phases, including hard and softmagnetic phases, exist. An average grain size of the hard magnetic phaseis equal to or greater than 10 nm and equal to or less than 200 nm,while that of the soft magnetic phase is equal to or greater than 1 nmand equal to or less than 100 nm.

[0036] In one embodiment of the present invention, the rapidlysolidified alloy made in the cooling step preferably includes an R₂Fe₁₄Bphase at 60 volume percent or more.

[0037] In another embodiment of the present invention, the cooling steppreferably includes rapidly cooling the melt within an ambient gas at apressure of 30 kPa or more to make a rapidly solidified alloy includingan R₂Fe₁₄B phase with an average grain size of 50 nm or less.

[0038] In this particular embodiment, the cooling step may include:bringing the melt into contact with the surface of a rotating chillroller to obtain a supercooled liquid alloy; and dissipating heat fromthe supercooled alloy into the ambient gas to grow the R₂Fe₁₄B phaseafter the supercooled alloy has left the chill roller.

[0039] In still another embodiment, the method may include the step ofheating and crystallizing the rapidly solidified alloy to form astructure in which three or more crystalline phases, including at leastR₂Fe₁₄B compound, α-Fe and boride phases, are included. In this processstep, an average crystal grain size of the R₂Fe₁₄B phase is set equal toor greater than 20 run and equal to or less than 150 nm, while that ofthe α-Fe and boride phases is set equal to or greater than 1 nm andequal to or less than 50 nm.

[0040] Specifically, the boride phase preferably includes an iron-basedboride with ferromagnetic properties.

[0041] More particularly, the iron-based boride preferably includes Fe₃Band/or Fe₂₃B₆.

[0042] In yet another embodiment, the melt may be cooled by a stripcasting process.

[0043] An inventive method for producing a bonded magnet includes thesteps of: preparing a powder of the iron-based rare earth alloy magnetby the second inventive method for producing the iron-based rare earthalloy magnet; and producing a bonded magnet using the powder of theiron-based rare earth alloy magnet.

BRIEF DESCRIPTION OF THE DRAWINGS

[0044]FIG. 1 is a graph illustrating a relationship between the maximumenergy product (BH)_(max) and the concentration of boron in an Nd—Fe—Btype nanocomposite magnet to which no Ti is added.

[0045]FIG. 2 is a graph illustrating a relationship between the maximumenergy product (BH)_(max) and the concentration of boron in an Nd—Fe—Btype nanocomposite magnet to which Ti is added.

[0046]FIG. 3 is a graph schematically illustrating secondary coolingeffects of an ambient gas by a relationship between the time passedsince a cooling process was started and the temperature of an alloy.

[0047]FIG. 4 schematically illustrates R₂Fe₁₄B compound and (Fe,Ti)—Bphases included in the magnet of the present invention.

[0048]FIG. 5 schematically illustrates how three types of rapidlysolidified alloys having a composition including additive Ti, acomposition including V or Cr as an alternative additive and acomposition including Nb, Mo or W as another alternative additive,respectively, change their microstructures during the crystallizationprocesses thereof.

[0049]FIG. 6A is a cross-sectional view illustrating an overallarrangement for a melt quenching machine for use to make a rapidlysolidified alloy for the iron-based rare earth alloy magnet of thepresent invention; and

[0050]FIG. 6B illustrates part of the machine shown in FIG. 6A, where amelt is quenched and rapidly solidified, to a larger scale.

[0051]FIG. 7 is a graph illustrating the demagnetization curves (i.e.,the second quadrant portion of the hysteresis loop) of samples Nos. 2and 3 and No. 11.

[0052]FIG. 8 is a graph illustrating the XRD patterns of theheat-treated samples Nos. 2, 3 and 11.

[0053]FIG. 9 is a graph illustrating the demagnetization curves ofsamples Nos. 13 and 17, respectively.

[0054]FIG. 10 is a graph illustrating the XRD (X-Ray Diffraction)patterns of the sample No. 13 before and after the sample is thermallytreated.

[0055]FIG. 11 is a graph illustrating the XRD patterns of the sample No.17 before and after the sample is thermally treated.

[0056]FIG. 12 is a graph illustrating the XRD patterns of sample No. 21before the sample is heated and crystallized and after it has beenheat-treated at 640° C. for 6 minutes.

[0057]FIG. 13 is a graph illustrating the demagnetization curves ofsamples Nos. 21 and 26, respectively.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0058] The iron-based rare earth alloy magnet of the present inventionis formed by rapidly cooling and solidifying a melt of a rare earthelement-iron-boron type alloy containing Ti. The rapidly solidifiedalloy includes crystalline phases. However, if necessary, the alloy isheated and further crystallized.

[0059] The present inventors found that when Ti is added to aniron-based rare earth alloy with a composition defined by a particularcombination of mole fraction ranges, the crystal growth of an α-Fephase, often observed while the melt is cooled, is suppressible, and thecrystal growth of an R₂Fe₁₄B phase can proceed preferentially anduniformly. The basic idea of the present invention lies in this finding.

[0060] Unless Ti is added to the material alloy, the α-Fe phase easilynucleates and grows faster and earlier than the crystal growth anNd₂Fe₁₄B phase. Accordingly, when the rapidly solidified alloy has beenthermally treated, the α-Fe phase with soft magnetic properties willhave grown excessively.

[0061] In contrast, where Ti is added to the material alloy, thecrystallization kinetics of the α-Fe phase would be slowed down, i.e.,it Would take a longer time for the a Fe phase to nucleate and grow.Thus, the present inventors believe the Nd₂Fe₁₄B phase would start tonucleate and grow before the α-Fe phase has grown coarsely. For thatreason, where Ti is added, crystal grains in the Nd₂Fe₁₄B phase can begrown sufficiently and dispersed uniformly before the α-Fe phase growstoo much. Furthermore, the addition of Ti appears to enhance thecrystallization of iron-based borides. Since Ti has a strong affinity toB, Ti stabilizes the iron-based borides by partitioning in the borides.

[0062] In the present invention, the additive Ti contributes toconsiderable reduction in grain size of the soft magnetic phases (e.g.,iron-based boride and α-Fe phases), uniform dispersion of the Nd₂Fe₁₄Bphase and increase in volume fraction of the Nd₂Fe₁₄B phase. As aresult, the composite magnet can have its coercivity and remanenceincreased sufficiently and can have the loop squareness of itsdemagnetization curve improved.

[0063] Hereinafter, the iron-based rare earth alloy magnet of thepresent invention will be described in further detail.

[0064] The inventive iron-based rare earth alloy magnet is preferablyrepresented by the general formula:(Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z) where T is at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare earth element substantially excluding La and Ce; and M is atleast one metal element selected from the group consisting of Ti, Zr andHf and always includes Ti.

[0065] The mole fractions x, y, z and m preferably meet the inequalitiesof: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1 at %≦z≦12 at %; and0≦m≦0.5, respectively.

[0066] The iron-based rare earth alloy magnet of the present inventionincludes a rare earth element at as small a mole fraction as 10 at % orless. However, since Ti has been added to the material alloy thereof,the inventive magnet can attain the unexpected effects of keeping, oreven increasing, the remanence B_(r) and improving the loop squarenessof the demagnetization curve thereof compared to an alloy not includingTi.

[0067] In the iron-based rare earth alloy magnet of the presentinvention, the soft magnetic phases have a very small grain size.Accordingly, the respective constituent phases are coupled togetherthrough exchange interactions. For that reason, even though softmagnetic phases, like iron-based boride and α-Fe phases, exist thereinin addition to the hard magnetic R₂Fe₁₄B phase, the alloy as a whole canshow excellent squareness at the demagnetization curve thereof.

[0068] The inventive iron-based rare earth alloy magnet preferablyincludes iron-based boride and α-Fe phases with a saturationmagnetization equal to, or even higher than, that of the R₂Fe₁₄B phase.Examples of the iron-based borides include Fe₃B (with a saturationmagnetization of 1.5 T) and Fe₂₃B₆ (with a saturation magnetization of1.6 T). In this case, the R₂Fe₁₄B phase has a saturation magnetizationof about 1.6 T and the α-Fe phase has a saturation magnetization of 2.1T.

[0069] Normally, where the mole fraction x of B is greater than 10 at %and the mole fraction y of the rare earth element R is from 6 at % to 8at %, R₂Fe₂₃B₃ is produced. However, even when a material alloy withsuch a composition is used, the addition of Ti can produce R₂Fe₁₄B,Fe₂₃B₆ and α-Fe, not R₂Fe₂₃B₃, in the present invention. Theseiron-based borides contribute to increasing the remanence.

[0070] As a result of experiments, the present inventors discovered thatonly when Ti was added, the remanence did not decrease but ratherincreased as opposed to any other metal element additive such as V, Cr,Mn, Nb or Mo. Also, when Ti was added, the loop squareness of thedemagnetization curve was much better than that obtained by adding anyof the elements cited above.

[0071] Furthermore, these effects attained by the additive Ti areparticularly remarkable where the concentration of B is greater than 10at %. Hereinafter, this point will be described with reference to FIGS.1 and 2.

[0072]FIG. 1 is a graph illustrating a relationship between the maximumenergy product (BH)_(max) and the concentration of boron in an Nd—Fe—Btype magnet to which no Ti is added. In FIG. 1, the white bars representdata about samples containing Nd at between 10 at % and 14 at %, whilethe black bars represent data about samples containing Nd at between 8at % and 10 at %. On the other hand, FIG. 2 is a graph illustrating arelationship between the maximum energy product (BH)_(max) and theconcentration of boron in an Nd—Fe—B type magnet to which Ti is added.In FIG. 2, the white bars also represent data about samples containingNd at between 10 at % and 14 at %, while the black bars also representdata about samples containing Nd at between 8 at % and 10 at %.

[0073] As can be seen from FIG. 1, once the concentration of boronexceeds 10 at %, the samples including no Ti have their maximum energyproducts (BH)_(max) decreased no matter how much Nd is containedtherein. Where the content of Nd is between 10 at % and 14 at %, thisdecrease is particularly noticeable. This tendency has been well knownin the art and it has been widely believed that any permanent magnet,including an Nd₂Fe₁₄B phase as its primary phase, should not containmore than 10 at % of boron. For instance, U.S. Pat. No. 4,836,868discloses a working example in which the concentration of boron is setto from 5 at % to 9.5 at %. This patent teaches that the concentrationof boron should preferably be equal to or greater than 4 at % and lessthan 12 at %, more preferably from 4 at % to 10 at %.

[0074] In contrast, as can be seen from FIG. 2, the samples includingthe additive Ti have their maximum energy products (BH)_(max) increasedin a certain range where the B concentration is greater than 10 at %.This increase is particularly remarkable where the Nd content is between8 at % and 10 at %.

[0075] Thus, the present invention can reverse the conventionalmisbelief that a B concentration of greater than 10 at % degrades themagnetic properties and can attain the unexpected effects just by addingTi to the material alloy.

[0076] Next, it will be described how to produce the iron-based rareearth alloy magnet of the present invention.

[0077] In the inventive process, a melt of the iron-based rare earthalloy with the above composition is rapidly cooled within an inert gasenvironment, thereby preparing a rapidly solidified alloy including anR₂Fe₁₄B phase at 60 volume % or more. The average grain size of theR₂Fe₁₄B phase in the quenched alloy is 50 nm or less, for example. Ifnecessary, this quenched alloy may be heat-treated. Then, the amorphousphases remaining in the alloy can be crystallized.

[0078] In a preferred embodiment, the melt is rapidly cooled within anenvironment at a pressure of 30 kPa or more. Then, the melt is not justrapidly cooled through the contact with a chill roller but also furthercooled due to secondary cooling effects caused by the ambient gas evenafter the solidified alloy left the roller.

[0079] By appropriately adjusting the circumference velocity of thechill roller, it is possible to transform the alloy into a supercooledliquid when the alloy leaves the chill roller. After leaving the chillroller, the alloy in the supercooled liquid state has its heatdissipated into the ambient gas so as to be crystallized.

[0080] Next, it will be described with reference to FIG. 3 how theambient gas brings about the secondary cooling effects. FIG. 3 is agraph schematically illustrating a relationship between the time passedsince the rapid cooling process was started and the temperature of thealloy. FIG. 3 illustrates two types of alloy cooling passages a and bassociated with relatively high and relatively low ambient gas pressuresof greater than 30 kPa and less than 30 kPa, respectively. The ranges inwhich α-Fe, Nd₂Fe₁₄B and Fe₂₃B₆ phases crystallize are illustrated inFIG. 3 as well, where T_(m) indicates the melting point of the alloy andT_(g) indicates the glass transition point of the alloy.

[0081] As can be seen from FIG. 3, where the ambient gas has arelatively low pressure (corresponding to the passage b), the ambientgas would not bring about so great secondary cooling effects.Accordingly, in that case, the surface velocity of the chill roller isset higher, thereby increasing the rapid cooling (i.e., primary cooling)rate of the chill roller. After leaving the surface of the chill roller,the alloy is slowly cooled by the ambient gas at a relatively low rate(i.e., the secondary cooling process). In FIG. 3, the node of thecooling passage b corresponds to a point in time the alloy leaves thechill roller.

[0082] On the other hand, where the ambient gas has a relatively highpressure (corresponding to the passage a), the ambient gas brings aboutremarkable secondary cooling effects, thus shortening the time it takesfor the alloy to pass the range where the Nd₂Fe₁₄B phase is produced.Probably for this reason, the growth of the Nd₂Fe₁₄B phase is suppressedand an Nd₂Fe₁₄B phase with a small grain size can be obtained instead.

[0083] As can be seen, where the pressure of the ambient gas is setlower than 30 kPa, the R₂Fe₁₄B phase, produced in the rapidly solidifiedalloy, will have an excessively large crystal grain size, therebydeteriorating the resultant magnet properties.

[0084] However, if the ambient gas has too high a pressure (i.e., higherthan the atmospheric pressure), then the ambient gas would be trappedbetween the melt and the chill roller and affects the cooling rateconsiderably. As a result, the chill roller cannot cool the meltsufficiently. Then, an α-Fe phase with a very large grain sizeprecipitates and therefore good hard magnetic properties cannot beattained.

[0085] According to the results of experiments the present inventorscarried out, while the rapid cooling process is performed, the ambientgas should have its pressure controlled preferably at 30 kPa or more andthe atmospheric pressure (=101.3 kPa) or less, more preferably between30 kPa and 90 kPa, and even more preferably between 40 kPa and 60 kPa.

[0086] Where the ambient gas has a pressure falling within any of thesepreferred ranges, the surface velocity of the chill roller is preferablyfrom 4 meters/second (=m/sec) to 50 n/sec. This is because if the rollersurface velocity is lower than 4 m/sec, then the R₂Fe₁₄B phase, includedin the rapidly solidified alloy, will have excessively large crystalgrains. In that case, the R₂Fe₁₄B phase will further increase its sizewhen thermally treated, thus possibly deteriorating the resultantmagnetic properties.

[0087] On the other hand, if the roller surface velocity is higher than50 m/sec, then the rapidly solidified alloy will be amorphized almostcompletely and substantially no R₂Fe₁₄B phase will precipitate.Accordingly, while the alloy is heated and crystallized, the R₂Fe₁₄Bphase will have its grain size increased considerably to have anon-uniform structure. As a result, the magnetic properties cannot beimproved.

[0088] According to the experimental results the present inventorsobtained, the roller surface velocity is more preferably from 5 m/sec to30 m/sec, even more preferably from 5 m/sec to 20 m/sec.

[0089] In the present invention, the rapidly solidified alloy has eithera structure in which almost no α-Fe phase with an excessively largegrain size precipitates but an R₂Fe₁₄B phase exists instead or astructure in which an R₂Fe₁₄B phase and an amorphous phase coexist.Accordingly, a high-performance composite magnet, in which crystalgrains in soft magnetic phases are dispersed finely or distributeduniformly in the grain boundary region between the hard magnetic phases,can be obtained. As used herein, the term “amorphous phase” means notonly a phase in which the atomic arrangement is sufficiently disordered,but also a phase containing embryos for crystallization, extremely smallcrystalline regions (size: several nanometers or less), and/or atomicclusters. More specifically, the term “amorphous phase” means the phasethat the crystal structure thereof cannot defined by an X-raydiffraction method or a TEM observation.

[0090] In the prior art, even when one tries to obtain a rapidlysolidified alloy including 60 volume % or more of R₂Fe₁₄B phase byrapidly cooling a molten alloy with a composition similar to that of thepresent invention, the resultant alloy will have a structure in which alot of α-Fe phase has grown coarsely. Thus, when the alloy is heated andcrystallized after that, the α-Fe phase will increase its grain sizeexcessively. Once soft magnetic phases, including the α-Fe phase, havegrown too much, the magnet properties of the alloy deterioratesignificantly, thus making it virtually impossible to produce a qualitypermanent magnet out of such an alloy.

[0091] Particularly with a material alloy containing boron at arelatively high percentage like the alloy of the present invention, evenif the melt is cooled at a low rate, crystalline phases cannot beproduced so easily according to the known method. This is because boronhighly likely creates an amorphous phase. For that reason, in the priorart, even if one tries to make a rapidly solidified alloy including 60volume % or more of R₂Fe₁₄B phase by decreasing the cooling rate of themelt sufficiently, not only the R₂Fe₁₄B phase but also the α-Fe phase orits precursor will precipitate a lot. Thus, when that alloy is heatedand crystallized after that, the α-Fe phase will further grow todeteriorate the magnet properties of the alloy seriously.

[0092] Thus, it was widely believed that the best way of obtaining ananocomposite magnet with a high coercivity is cooling a melt at anincreased rate to amorphize most of it first and then forming a highlyfine and uniform structure by heating the amorphous phases. This isbecause they took it for granted that there should be no otheralternative but crystallizing the amorphous phases through an easilycontrollable heat treatment process to obtain a nanocomposite structurein which fine crystal grains are dispersed finely.

[0093] On this popular belief, W. C. Chan et al., reported a techniqueof obtaining Nd₂Fe₁₄B and α-Fe phases with grain sizes on the order ofseveral tens nm. According to Chan's technique, La, which excels inproducing the amorphous phases, is added to a material alloy. Next, thematerial alloy is melt quenched to obtain a rapidly solidified alloymainly composed of the amorphous phases. And then the alloy is heatedand crystallized. See W. C. Chan et al., “The Effects of RefractoryMetals on the Magnetic Properties of α-Fe/R₂Fe₁₄B-type Nanocomposites”,IEEE Trans. Magn. No.5, INTER-MAG. 99, Kyongiu, Korea, pp.3265-3267,1999. This article also teaches that adding a refractory metal elementsuch as Ti in a very small amount (e.g., 2 at %) improves the magneticproperties and that the mole fraction of Nd, rare earth element, ispreferably increased from 9.5 at % to 11.0 at % to reduce the grainsizes of the Nd₂Fe₁₄B and α-Fe phases. The refractory metal is added tosuppress borides like R₂Fe₂₃B₃ and Fe₃B from being produced and to makea magnet consisting of Nd₂Fe₁₄B and α-Fe phases only.

[0094] In contrast, according to the present invention, Ti added cansuppress not only the nucleation of the α-Fe phase during the rapidsolidification process but also the grain growth of soft magnetic phaseslike iron-based boride and α-Fe phases during the heating/crystallizingprocess.

[0095] According to the present invention, even though the materialalloy used includes a rare earth element at a relatively low percentage(i.e., 9.3 at % or less), a permanent magnet, exhibiting high remanenceand coercivity and showing excellent loop squareness in itsdemagnetization curve, can be obtained.

[0096] As described above, the inventive magnet can have its coercivityincreased by making the Nd₂Fe₁₄B phase nucleate and grow faster andearlier in the cooling process so that the Nd₂Fe₁₄B phase increases itsvolume percentage and yet by suppressing the grain coarsening of thesoft magnetic phases. Also, the remanence thereof increases probablybecause Ti added can produce a boride phase (e.g., ferromagneticiron-based borides) from the B-rich non-magnetic amorphous phasesexisting in the rapidly solidified alloy and can reduce the volumepercentage of the non-magnetic amorphous phases remaining in the heatedand crystallized alloy.

[0097] The rapidly solidified alloy obtained this way is preferablyheated and crystallized if necessary to form a structure with three ormore crystalline phases including R₂Fe₁₄B compound, boride and α-Fephases. The heat treatment is preferably conducted with its temperatureand duration controlled in such a manner that the R₂Fe₁₄B phase willhave an average crystal grain size of 10 nm to 200 nm and that theboride and α-Fe phases will have an average crystal grain size of 1 nmto 50 nm. The R₂Fe₁₄B phase normally has an average crystal grain sizeof 30 nm or more, which may be 50 nm or more depending on theconditions. On the other hand, the soft magnetic phases, like boride andα-Fe phases, often have an average crystal grain size of 30 nm or lessand typically several nanometers at most.

[0098] In the resultant magnet, the R₂Fe₁₄B phase has an average crystalgrain size greater than that of the α-Fe phase. FIG. 4 schematicallyillustrates the metal structure of this magnet. As shown in FIG. 4, finecrystal grains in soft magnetic phases are distributed in a grainboundary of relatively large crystal grains in the R₂Fe₁₄B phase. Eventhough the R₂Fe₁₄B phase has a relatively large average grain size, thesoft magnetic phases have a relatively small average grain size.Accordingly, these constituent phases are coupled together throughexchange interaction and the magnetization direction of the softmagnetic phases is constrained by the hard magnetic phase. Consequently,the alloy as a whole can show excellent loop squareness in itsdemagnetization curve.

[0099] In the inventive process, borides are easily produced asdescribed above. The reason is probably as follows. When a rapidlysolidified alloy, mostly composed of the R₂Fe₁₄B phase, is made, theamorphous phases existing in the solidified alloy should contain anexcessive amount of boron. Accordingly, when the alloy is heated andcrystallized, that boron will bond to other elements easily, thusnucleating and growing in profusion. However, if that boron, containedin the amorphous phases before the heat treatment, bonds to otherelements and produces compounds with low remanences, then the magnet asa whole will have its remanence decreased.

[0100] As a result of experiments, the present inventors discovered thatonly when Ti was added, the remanence did not decrease but ratherincreased as opposed to any other metal element additive such as V, Cr,Mn, Nb or Mo. Also, when M (e.g., Ti, in particular) was added, the loopsquareness of the demagnetization curve was much better than any of theelements cited above did.

[0101] Accordingly, the present inventors believe that Ti plays a keyrole in suppressing the production of borides with low remanences.Particularly when relatively small amounts of B and Ti are included in amaterial alloy with the composition defined by the present invention,iron-based borides with ferromagnetic properties will easily grow duringthe alloy is heat-treated. In such a case, boron included in thenon-magnetic amorphous phases would be absorbed into the iron-basedborides. For that reason, the nonmagnetic amorphous phases, remaining inthe alloy after the alloy has been heated and crystallized, decreasetheir volume percentage but the ferromagnetic crystalline phaseincreases its volume percentage instead, thus increasing the remanenceBr.

[0102] Hereinafter, this point will be further detailed with referenceto FIG. 5.

[0103]FIG. 5 schematically illustrates how three types of rapidlysolidified alloys having a composition including additive Ti, acomposition including V or Cr as an alternative additive and acomposition including Nb, Mo or W as another alternative additive,respectively, change their microstructures during the crystallizationprocesses thereof. Where Ti is added, the grain growth of the softmagnetic phases is suppressed in a temperature range exceeding thetemperature at which the α-Fe phase grows rapidly. As a result,excellent hard magnetic properties can be maintained. In contrast, whereany of the other metal elements (e.g., Nb, V, Cr, etc.) is added, thegrain growth of the respective constituent phases advances remarkablyand the exchange coupling among those phases weakens in the relativelyhigh temperature range in which the α-Fe phase grows rapidly. As aresult, the resultant demagnetization curves have their loop squarenessdecreased.

[0104] First, a situation where Nb, Mo or W is added will be described.In this case, if the alloy is thermally treated in a relatively lowtemperature range where no α-Fe phase precipitates, then good hardmagnetic properties, including superior loop squareness of thedemagnetization curve, are attainable. In an alloy that was heat-treatedat such a low temperature, however, R₂Fe₁₄B crystalline particles wouldbe dispersed in the non-magnetic amorphous phases and the alloy does nothave the nanocomposite magnet structure. Also, if the alloy isheat-treated at an even higher temperature, then the α-Fe phasenucleates and grows out of the amorphous phases. Unlike the situationwhere Ti is added, the α-Fe phase rapidly grows and increases its grainsize excessively. As a result, the exchange coupling among theconstituent phases weakens and the loop squareness of thedemagnetization curve deteriorates considerably.

[0105] On the other hand, where Ti is added, a nanocomposite structure,including microcrystalline R₂Fe₁₄B, iron-based boride, and α-Fe can beobtained by heat-treating the alloy and the grains in the respectiveconstituent phases are dispersed finely and uniformly. Also, theaddition of Ti suppresses the grain growth of the α-Fe phase.

[0106] Where V or Cr is added, any of these additive metal elements iscoupled anti-ferromagnetically with Fe to form a solid solution, thusdecreasing the remanence considerably. The additive V or Cr cannotsuppress the heat-treatment-induced grain growth sufficiently, either,and deteriorates the loop squareness of the demagnetization curve.

[0107] Accordingly, only when Ti is added, the crystal grain coarseningof the α-Fe phase can be suppressed appropriately and iron-based borideswith ferromagnetic properties can be obtained. Furthermore, Ti, as wellas B and C, plays an important role as an element that delays thecrystallization of Fe initial crystals (i.e., γ-Fe that will betransformed into α-Fe) during the rapid cooling process and therebyfacilitates the production of the supercooled liquid. Accordingly, evenif the melt of the alloy including Ti is rapidly cooled and solidifiedat a relatively low cooling rate between about 10²° C./sec and about10⁵° C./sec, a rapidly solidified alloy, in which the α-Fe phase has notprecipitated too much and the microcrystalline R₂Fe₁₄B and amorphousphases coexist, can be obtained. This greatly contributes to the costreduction of nanocomposite magnets because this means that a stripcasting method, particularly suitable for mass production, can beemployed as a rapid cooling technique.

[0108] The strip casting technique is a highly productive andcost-effective method for obtaining a solidified alloy by rapidlycooling a melt of a material alloy. This is because in the strip castingmethod, the flow rate of the melt does not have to be controlled usingthe nozzle orifice but the melt may be poured directly from a turndishonto a chill roller. To amorphize the melt of an R—Fe—B rare earth alloyin a cooling rate range applicable to even the strip casting method,normally B (boron) should be added at 10 at % or more. However, if Badded is that much, then not just non-magnetic amorphous phases but alsoα-Fe phase and/or soft magnetic Nd₂Fe₂₃B₃ phase will growpreferentially. That is to say, no uniform, microcrystalline structurecan be obtained. As a result, the volume percentage of the ferromagneticphases decreases, the remanence drops and the volume percentage of theNd₂Fe₁₄B phase also decreases. Consequently, the coercivity decreasesnoticeably. However, if Ti is added as in the present invention, thenthe remanence increases.

[0109] It should be noted that a rapidly solidified alloy, including theNd₂Fe₁₄B phase at a high volume percentage, could improve the resultantmagnet properties more easily than a solidified alloy including theamorphous phases at a high volume percentage. Accordingly, the volumepercentage of the Nd₂Fe₁₄B phase to the entire solidified alloy ispreferably 50 volume % or more, more specifically 60 volume % or more,which value was obtained by Mossbauer spectroscopy.

Preferred Composition

[0110] Q is either B (boron) only or a combination of B and C (carbon).The molar fraction of C to Q is preferably 0.25 or less.

[0111] If the mole fraction x of Q is 10 at % or less, then it isdifficult to make the desired rapidly solidified alloy, in which themicrocrystalline R₂Fe₁₄B and amorphous phases coexist, at the lowcooling rate between about 10²° C./sec to about 10⁵° C./sec. Also, evenif the alloy is heat-treated after that, the resultant H_(cJ) will be aslow as less than 480 kA/m. In addition, the strip casting method, whichis one of the most cost-effective techniques among various rapid coolingmethods, cannot be adopted in that case, and the price of the resultantpermanent magnet product rises unintentionally. On the other hand, ifthe mole fraction x of Q exceeds 20 at %, then the volume percentage ofthe amorphous phases, remaining in the alloy even after the alloy hasbeen heated and crystallized, increases. In addition, the percentage ofthe α-Fe phase, which has a higher saturation magnetization than anyother constituent phase, decreases and the remanence B_(r) drops. Inview of these respects, the mole fraction x of Q is preferably greaterthan 10 at % and equal to or less than 20 at %, more preferably greaterthan 10 at % and less than 17 at/o.

[0112] R is at least one element selected from the rare earth elementsand/or yttrium (Y). Preferably, R includes substantially no La andsubstantially no Ce, because the existence of La or Ce decreases thecoercivity and the loop squareness of the demanetization curve. However,there is no problem of degrading the magnetic properties if very smallamounts (i.e., 0.5 at % or less) of La and Ce exist as inevitableimpurities. Therefore, the term “substantially no La (Ce)” or“substantially excluding La (Ce)” means that the content of La (Ce) is0.5 at % or less. R preferably includes Pr or Nd as an indispensableelement, part of which may be replaced with Dy and/or Th. If the molefraction y of R is less than 6 at %, then fine crystal grains with themicrocrystalline R₂Fe₁₄B structure, which is needed for expressing thecoercivity, do not crystallize sufficiently and a coercivity H_(cJ) of480 kA/m or more cannot be obtained. On the other hand, if the molefraction y of R is equal to or greater than 10 at %, then thepercentages of the iron-based borides and a Fe with ferromagneticproperties both decrease. For these reasons, the mole fraction y of therare earth element R is preferably equal to or greater than 6 at % andless than 10 at % (e.g., from 6 at % to 9.5 at %), more preferably 8 at% to 9.3 at %, and most preferably from 8.3 at % to 9.0 at %.

[0113] The additive element(s) M must include Ti and may further includeZr and/or Hf optionally. To attain the above effects, Ti isindispensable. As described above, the additive Ti increases thecoercivity H_(cJ), remanence B_(r) and maximum energy product (BH)_(max)and improves the loop squareness of the demagnetization curve.

[0114] If the mole fraction z of the metal element(s) M is less than 0.5at %, then the above effects cannot be attained fully even though Ti isadded. Nevertheless, if the mole fraction z of the metal element(s) Mexceeds 12 at %, then the volume percentage of the amorphous phases,remaining in the alloy even after the alloy has been heated andcrystallized, increases and the remanence B_(r) likely drops. In view ofthese respects the mole fraction z of the metal element(s) M ispreferably from 0.5 at % to 12 at %. The lower limit of a morepreferable z range is 1.0 at % and the upper limit thereof is 8.0 at %.The upper limit of an even more preferable z range is 6.0 at %.

[0115] Also, the higher the mole fraction x of Q, the more likely theamorphous phases, including an excessive amount of Q (e.g., boron), areformed. Accordingly, the mole fraction z of the metal element(s) Mshould preferably be set higher because of this reason also.Specifically, the mole fractions x and z should be adjusted preferablyto meet the inequality z/x≧0.1, more preferably to meet the inequalityz/x≧0.15.

[0116] As described above, the metal element(s) M must include Tibecause Ti acts very favorably. In this case, the atomic percentage ofTi to the total metal element(s) M is preferably 70 at % or more andmore preferably 90 at % or more.

[0117] The balance of the alloy, other than the elements Q, R and M, maybe Fe alone. Or, part of Fe may be replaced with at least one transitionmetal element T selected from the group consisting of Co and Ni, becausethe desired hard magnetic properties are also attainable. However, ifmore than 50% of Fe is replaced with T, then a high remanence B_(r) of0.7 T or more cannot be obtained. For that reason, the percentage of Fereplaced is preferably from 0% to 50%. Also, by replacing part of Fewith Co, the loop squareness of the demagnetization curve improves andthe Curie temperature of the R₂Fe₁₄B phase rises, thus increasing thethermal resistance of the alloy. The percentage of Fe replaced with Cois preferably from 0.5% to 40%.

[0118] Hereinafter, preferred embodiments of the present invention willbe described with reference to the accompanying drawings.

Melt Quenching Machine

[0119] In this embodiment, a material alloy is prepared using a meltquenching machine such as that shown in FIGS. 6A and 6B. The alloypreparation process is performed within an inert gas environment toprevent the material alloy, which contains rare earth element R and Fethat are easily oxidizable, from being oxidized. The inert gas may beeither a rare gas of helium or argon, for example or nitrogen. The raregas of helium or argon is preferred to nitrogen, because nitrogen reactswith the rare earth element R relatively easily.

[0120] The machine shown in FIG. 6A includes material alloy melting andquenching chambers 1 and 2, in which a vacuum or an inert gasenvironment is created at an adjustable pressure. Specifically, FIG. 6Aillustrates an overall arrangement for the machine, while FIG. 6Billustrates part of the machine to a larger scale.

[0121] As shown in FIG. 6A, the melting chamber 1 includes a meltingfurnace 3, a melt crucible 4 with a teeming nozzle 5 at the bottom andan airtight compounded material feeder 8. A material alloy 20, which hasbeen compounded to have a desired magnet alloy composition and suppliedfrom the feeder 8, is melted in the melting furnace 3 at an elevatedtemperature. A melt 21 of the material alloy 20 is poured into thecrucible 4, which is provided with a heater (not shown) for keeping thetemperature of the melt teemed therefrom at a predetermined level.

[0122] The quenching chamber 2 includes a rotating chill roller 7 forrapidly cooling and solidifying the melt 21 that has been drippedthrough the teeming nozzle 5.

[0123] In this machine, the environment and pressure inside the meltingand quenching chambers 1 and 2 are controllable within prescribedranges. For that purpose, ambient gas inlet ports 1 b, 2 b and 8 b andoutlet ports 1 a, 2 a and 8 a are provided at appropriate positions ofthe machine. In particular, the gas outlet port 2 a is connected to apump to control the absolute pressure inside the quenching chamber 2within a range from 30 kPa to atmospheric pressure.

[0124] The melting furnace 3 may be inclined at a desired angle to pourthe melt 21 through a funnel 6 into the crucible 4. The melt 21 isheated in the crucible 4 by the heater (not shown).

[0125] The teeming nozzle 5 of the crucible 4 is positioned on theboundary wall between the melting and quenching chambers 1 and 2 to dripthe melt 21 in the crucible 4 onto the surface of the chill roller 7,which is located under the nozzle 5. The orifice diameter of the nozzle5 may be in a range from 0.5 mm to 2.0 mm, for example. If the viscosityof the melt 21 is high, then the melt 21 cannot flow through the nozzle5 easily. In this embodiment, however, the pressure inside the quenchingchamber 2 is kept lower than the pressure inside the melting chamber 1.Accordingly, there exists an appropriate pressure difference between themelting and quenching chambers 1 and 2, and the melt 21 can be teemedsmoothly.

[0126] To attain a good thermal conductivity, the chill roler 7 may bemade of Al alloy, Cu alloy, carbon steel, brass, W, Mo or bronze.However, the roller 7 is preferably made of Cu, because Cu realizes asufficient mechanical strength at a reasonable cost. The diameter of theroller 7 may be in a range from 300 mm to 500 mm, for instance. Thewater-cooling capability of a water cooler provided inside the roller 7is calculated and adjustable based on the latent heat of solidificationand the volume of the melt teemed per unit time.

[0127] The machine shown in FIGS. 6A and 6B can rapidly solidify 10 kgof material alloy in 10 to 20 minutes, for example. The alloy solidifiedin this manner is in the form of a thin strip (or ribbon) 22 with athickness of 100 μm to 300 μm and a width of 2 mm to 3 mm.

Melt Quenching Process

[0128] First, the melt 21 of the material alloy, which is represented bythe general formula described above, is prepared and stored in thecrucible 4 of the melting chamber 1 shown in FIG. 6A. Next, the melt 21is dripped through the teeming nozzle 5 onto the chill roller 7 to comeinto contact with, and be rapidly cooled and solidified by, the roller 7within a low-pressure Ar environment. In this case, an appropriate rapidsolidification technique, making the cooling rate controllableprecisely, should be adopted.

[0129] In the illustrated embodiment, the melt 21 should be quenched andsolidified preferably at a rate between 1×10²° C./sec and 1×10⁸° C./sec,more preferably at a rate between 1×10⁴° C./sec and 1×10⁶° C./sec.

[0130] An interval during which the melt 21 is quenched by the chillroller 7 is equivalent to an interval between a point in time the alloycomes into contact with the circumference of the rotating chill roller 7and a point in time the alloy leaves the roller 7. In the meantime, thealloy has its temperature decreased to become a supercooled liquid.Thereafter, the supercooled alloy leaves the roller 7 and travels withinthe inert gas environment. While the thin-strip alloy is travelling, thealloy has its heat dissipated into the ambient gas. As a result, thetemperature of the alloy further drops. According to the presentinvention, the pressure of the ambient gas is in the range from 30 kPato atmospheric pressure. Thus, the heat of the alloy can be dissipatedinto the ambient gas even more effectively and the Nd₂Fe₁₄B phase cannucleate and grow uniformly in the alloy. It should be noted that unlessan appropriate amount of element M (e.g., Ti) has been added to thematerial alloy, then the α-Fe phase nucleates and grows faster andearlier in the rapidly solidified alloy, thus deteriorating theresultant magnet properties.

[0131] In this embodiment, the surface velocity of the roller 7 isadjusted to fall within a range from 10 m/sec to 30 m/sec and thepressure of the ambient gas is set to 30 kPa or more to enhance thesecondary cooling effects caused by the ambient gas. In this manner, arapidly solidified alloy, including 60 volume percent or more of R₂Fe₁₄Bphase with an average grain size of as small as 80 nm or less, isprepared.

[0132] In the present invention, the technique of rapidly cooling themelt is not limited to the single roller melt-spinning method describedabove. Examples of other applicable techniques include twin rollermethod, gas atomization method, strip casting method requiring no flowrate control using nozzle or orifice, and rapid cooling techniqueutilizing the roller and atomization methods in combination.

[0133] Among these rapid cooling techniques, the strip casting methodresults in a relatively low cooling rate, i.e., 10²° C./sec to 10⁵°C./sec. According to this embodiment, by adding an appropriate volume ofTi to the material alloy, a rapidly solidified alloy, most of which hasa structure including no Fe initial crystals, can be obtained even bythe strip casting method. The process cost of the strip casting methodcan be about half or less of any other rapid cooling method.Accordingly, in preparing a large quantity of rapidly solidified alloy,the strip casting method is much more effective than the single rollermethod, and is suitably applicable to mass production. However, if noelement M is added to the material alloy or if Cr, V, Mn, Mo, Ta and/orW are/is added thereto instead of Ti, then a metal structure including alot of Fe initial crystals will be produced even in the rapidlysolidified alloy prepared by the strip casting method. Consequently, thedesired metal structure cannot be obtained.

Heat Treatment

[0134] In this embodiment, the heat treatment is conducted within anargon environment. Preferably, the alloy is heated at a temperature riserate of 5° C./sec to 20° C./sec, kept at a temperature between 550° C.and 850° C. for a period of time from 30 seconds to 20 minutes and thencooled to room temperature. This heat treatment results in nucleationand/or crystal growth of metastable phases in a remaining amorphousphase, thus forming a nanocomposite microcrystalline structure.According to the present invention, the microcrystalline Nd₂Fe₁₄B phasealready accounts for 60 volume % or more of the total alloy when theheat treatment is started. Thus, when the heat treatment is conductedunder these conditions, soft magnetic phases will not increase theirsize too much and the soft magnetic phases will be dispersed finely anduniformly in a grain boundary between the microcrystalline Nd₂Fe₁₄Bgrains.

[0135] If the heat treatment temperature is lower than 550° C., then alot of amorphous phases may remain even after the heat treatment and theresultant coercivity may not reach the desired level depending on theconditions of the rapid cooling process. On the other hand, if the heattreatment temperature exceeds 850° C., the grain growth of therespective constituent phases will advance too much, thus decreasing theremanence B_(r) and deteriorating the loop squareness of thedemagnetization curve. For these reasons, the heat treatment temperatureis preferably from 550° C. to 850° C., more preferably 570° C. to 820°C.

[0136] In the present invention, the ambient gas causes the secondarycooling effects so that a sufficient amount of crystal grains in theNd₂Fe₁₄B phase crystallize uniformly and finely in the rapidlysolidified alloy. Accordingly, even if the rapidly solidified alloy isnot heat-treated, the solidified alloy itself can exhibit good enoughmagnet properties. That is to say, the heat treatment forcrystallization is not indispensable for the present invention. However,to further improve the magnet properties, the heat treatment ispreferably conducted. In addition, even though the heat treatment iscarried out at lower temperatures than the known process, the magnetproperties still can be improved sufficiently.

[0137] To prevent the alloy from being oxidized, the heat treatment ispreferably conducted within an inert gas (e.g., Ar or N₂ gas)environment. The heat treatment may also be performed within a vacuum of0.1 kPa or less.

[0138] Before the heat treatment, the rapidly solidified alloy mayinclude metastable phases such as Fe₃B, Fe₂₃B₆, and R₂Fe₂₃B₃ phases inaddition to the R₂Fe₁₄B compound and amorphous phases. In that case,when the heat treatment is over, the R₂Fe₂₃B₃ phase will havedisappeared. Instead, crystal grains of an iron-based boride (e.g.,Fe₂₃B₆), showing a saturation magnetization equal to or even higher thanthat of R₂Fe₁₄B phase, or α-Fe phase can be grown.

[0139] In the present invention, even if the soft magnet phases like theα-Fe phase exist in the resultant magnet, excellent magnetic propertiesstill can be attained because the soft and hard magnetic phases aremagnetically coupled together through exchange interaction.

[0140] After the heat treatment, the Nd₂Fe₁₄B phase should have anaverage crystal grain size of 300 nm or less, which is a single magneticdomain size. The average crystal grain size of the Nd₂Fe₁₄B phase ispreferably from 20 nm to 150 nm, more preferably 20 nm to 100 nm. On theother hand, if the boride and α-Fe phases have an average crystal grainsize of more than 50 nm, then the exchange interaction among therespective constituent phases weakens, thus deteriorating the loopsquareness of the demagnetization curve and decreasing (BH)_(max). Andif the average crystal grain size of these phases is less than 1 nm,then a high coercivity cannot be attained. For these reasons, the softmagnet phases, such as the boride and α-Fe phases, should preferablyhave an average crystal grain size of 1 nm to 50 nm, more preferably 30nm or less.

[0141] It should be noted that the thin strip of the rapidly solidifiedalloy may be roughly cut or pulverized before subjected to the heattreatment.

[0142] After heat-treated, the resultant magnetic alloy is finelypulverized to obtain a magnet powder. Then, various types of bondedmagnets can be made from this powder by performing known process stepson this powder. In making a bonded magnet, the magnet powder of theiron-based rare earth alloy is compounded with an epoxy or nylon resinbinder and then molded into a desired shape. At this time, a magnetpowder of any other type (e.g., an Sm—Fe—N magnet powder or hard ferritemagnet powder) may be mixed with the nanocomposite magnet powder.

[0143] Using the resultant bonded magnet, motors, actuators and otherrotating machines can be produced.

[0144] When the magnet powder of the present invention is used for aninjection-molded bonded magnet, the powder is preferably pulverized tohave a mean particle size of 200 μm or less, more preferably from 30 μmto 150 μm. On the other hand, where the inventive magnet powder is usedfor a compression-molded bonded magnet, the powder is preferablypulverized to have a mean particle size of 300 μm or less, morepreferably from 30 μm to 250 μm and even more preferably 50 μm to 200 μmwith a bimodal size distribution.

[0145] It should be noted that if the powder obtained is subjected to asurface treatment (e.g., coupling treatment, conversion coating orplating), then the powder for a bonded magnet can have its moldabilityimproved no matter how the powder is molded. In addition, the resultantbonded magnet can have its anticorrosiveness and thermal resistance bothincreased. Alternatively, after a bonded magnet has been once formed bymolding the powder into a desired shape, the surface of the magnet mayalso be treated, e.g., covered with a plastic or conversion coating orplated. This is because anticorrosiveness and thermal resistance of thebonded magnet can also be increased in that case.

EXAMPLES

[0146] First, examples and comparative examples, in which the molefractions x and z of Q and M meet 10 at %<x<15 at % and 0.1 at %<z<10 at%, respectively, will be described.

[0147] For each of samples Nos. 1 to 12 shown in the following Table 1,the respective materials B, C, Fe, Co, Ti, Nd, Pr, Th and Dy withpurities of 99.5% or more were weighed so that the sample had a totalweight of 30 g and then the mixture was injected into a crucible ofquartz. In Table 1, samples Nos. 1 to 8 represent examples of thepresent invention, while samples Nos. 9 to 12 represent comparativeexamples: Heat Roller Treatment Composition (at %) Velocity temperatureFe Q R M m/sec ° C. 1 Fe79 B11 Nd9 Ti1 20.0 660 2 Fe78.7 B10.3 Nd9 Ti212.0 700 3 Fe76.7 B10.3 Nd9 Ti4 9.0 760 4 Fe69 + B14 Nd3 + Pr3 Ti6 9.0740 Co3 5 Fe68 + B7 + C4 Nd9.5 Ti8 7.0 780 Co3.5 6 Fe78.7 B10.3 Nd8 +Dy1 Ti2 12.0 720 7 Fe78.7 B5 + C5.3 Nd8 + Tb1 Ti2 12.0 720 8 Fe65.7 +B10.3 Nd9 Ti5 8.0 720 Co10 9 Fe81 B12 Nd7 — 30.0 660 10 Fe80 B14 Nd6 —20.0 680 11 Fe80.7 B10.3 Nd9 — 25.0 660 12 Fe76.7 B10.3 Nd11 Ti2 12.0710

[0148] In Table 1, the column “Q” includes “B7+C4”, for example, whichmeans that 7 at % of boron (B) and 4 at % of carbon (C) were added.Also, the column “R” includes “Nd3.+Pr3”, for example, which means that3 at % of Nd and 3 at % of Pr were added.

[0149] The quartz crucible used for preparing the melt had an orificewith a diameter of 0.8 mm at the bottom. Accordingly, the alloyincluding these materials was melted in the quartz crucible to be amelt, which was then dripped down through the orifice. The materialalloy was melted by a high frequency heating method within an argonenvironment at a pressure of 1.33 kPa. In the illustrated examples, thetemperature of the melt was set to 1500° C.

[0150] The surface of the melt was pressurized with Ar gas at 26.7 kPa,thereby propelling the melt against the outer circumference of a copperchill roller, which was located 0.7 mm under the orifice. The roller wasrotated at a high velocity while being cooled inside so that the outercircumference would have its temperature kept at around roomtemperature. Accordingly, the melt, which had been dripped down throughthe orifice, came into contact with the surface of the chill roller tohave its heat dissipated therefrom while being forced to rapidly move onthe rotating chill roller. The melt was continuously expelled throughthe orifice onto the surface of the roller. Thus, the rapidly solidifiedalloy was in the shape of an elongated thin strip (or ribbon) with awidth of 2 to 3 mm and a thickness of 20 to 50 μm.

[0151] In the (single) roller method adopted in these examples, thecooling rate is determined by the circumference velocity of the rollerand the weight of the melt dripped per unit time, which depends on thediameter (or cross-sectional area) of the orifice and the pressure onthe melt. In the present examples, the orifice has a diameter of 0.8 mm,a pressure of 26.7 kPa was placed on the melt and the dripping rate wasfrom about 0.5 kg/min to 1 kg/min.

[0152] The circumference velocity of the roller was changed as shown inTable 1.

[0153] Next, the rapidly solidified alloy samples Nos. 1 to 12 wereheat-treated within Ar gas. Specifically, the rapidly solidified alloyswere kept at the respective heat treatment temperatures shown on therightmost column in Table 1 for 6 minutes and then cooled to roomtemperature. Thereafter, the magnetic properties of these samples weremeasured using a vibrating sample magnetometer. The results are shown inthe following Table 2: TABLE 2 Magnetic properties B_(r) H_(CJ)(BH)_(max) (T) (kA/m) (kJ/m³) 1 0.86 490 94 2 0.85 605 118 3 0.85 695111 4 0.88 520 102 5 0.84 740 106 6 0.84 658 101 7 0.83 682 98 8 0.87730 125 9 0.80 22 4 10 0.69 8 — 11 0.86 479 80 12 0.74 955 88

[0154] As can be seen from Table 2, the examples of the presentinvention showed magnetic properties that were superior to those of thecomparative examples. Also, even if Ti had been added, its beneficialeffects (i.e., a highly uniform distribution of fine crystal grains)could not be attained fully and the remanence B_(r) dropped noticeablywhere the mole fraction y of the rare earth element R (Nd) was out ofthe range 6 at %≦y<10 at %.

[0155]FIG. 7 illustrates the demagnetization curves of samples Nos. 2and 3 (examples of the present invention) and sample No. 11 (comparativeexample). In FIG. 7, the ordinate represents the remanence while theabscissa represents the demagnetization field strength.

[0156] As can be seen from FIG. 7, the demagnetization curves of theexamples have loop squareness much better than that of the comparativeexample. The comparative example showed deteriorated loop squarenessprobably because the crystal grain size was large.

[0157] Next, each of samples Nos. 1 to 8 representing the examples ofthe present invention was analyzed using Cu—Kα characteristicX-radiation to identify the respective constituent phases thereof. As aresult, the existence of Fe₂₃B₆ and α-Fe phases, as well as R₂Fe₁₄Bphase, was confirmed. As for samples Nos. 9 and 10 representing two ofthe comparative examples on the other hand, no R₂Fe₁₄B phase, which is ahard magnetic phase, was identified and it was found that those samplesconsisted of R₂Fe₂₃B3 and α-Fe phases, which are soft magnetic phases.In sample No. 11 representing another comparative example, a mixture ofhard magnetic R₂Fe₁₄B and soft magnetic α-Fe phases was identified butno ferromagnetic iron-based borides were found.

[0158]FIG. 8 illustrates the XRD patterns of the heat-treated samplesNos. 2 and 3 (examples of the present invention) and No. 11 (comparativeexample). In FIG. 8, the ordinate represents the intensity ofdiffraction, while the abscissa represents the angle of diffraction.

[0159] As can be seen from FIG. 8, it was found that a metallicstructure, comprised of Nd₂Fe₁₄B, α-Fe and Fe₂₃B₆ phases, had beenformed in the examples of the present invention. As for the comparativeexample on the other hand, only Nd₂Fe₁₄B and α-Fe phases wereidentified. Accordingly, it is believed that an excessive amount of Bexisted in the resultant alloy structure.

[0160] Using a transmission electron microscope (TEM), the metalstructure after the heat treatment was investigated on samples Nos. 1 to8 representing the examples of the present invention. As a result, thepresent inventors found that each of these samples had ananometer-scaled crystalline structure with an average crystal grainsize of 10 to 25 nm. The present inventors also analyzed sample No. 2using an atom probe to find that part of Ti had been replaced with Fecontained in the respective constituent phases but that most of Tiexisted in the grain boundary phases

[0161] Next, other examples of the present invention and referenceexamples will be described. In each of the examples of the presentinvention, the mole fractions x and z of Q and M meet 15 at %≦x≦20 at %and 3.0 at %<z<12 at %, respectively. On the other hand, in thereference examples, the mole fraction x of Q does not meet 15 at %≦x≦=20at %, nor does the mole fraction z of M meet 3.0 at %<z<12 at %.

[0162] For each of samples Nos. 13 to 19 shown in the following Table 3,the respective materials B, C, Fe, Co, Ti and Nd with purities of 99.5%or more were weighed so that the sample had a total weight of 30 g andthen the mixture was injected into a crucible of quartz. TABLE 3 RollerHeat Surface treatment Composition (at %) velocity temperature Fe Q R Mm/sec ° C. 13 Fe68.5 B15 Nd8.5 Ti8 20 680 14 Fe70.0 + Co2.5 B15 Nd8.5Ti4 20 680 15 Fe71.5 B14 + C1 Nd8.5 Ti5 12 700 16 Fe66.5 B15 Nd8.5 Ti1025 720 17 Fe76.5 B15 Nd8.5 — 30 760 18 Fe74.5 B15 Nd8.5 Ti2 15 780 19Fe75.5 B15 Nd6.5 Ti3 20 780

[0163] In Table 3, the column “M” includes “Ti8”, for example, whichmeans that 8 at % of Ti was added to the material alloy, and “-” meansthat no Ti was added thereto.

[0164] Samples Nos. 13 to 19 were subjected to a rapid solidificationprocess under the same conditions as those specified for samples Nos. 1to 12.

[0165] The resultant rapidly solidified alloy structures were analyzedusing Cu—Kα characteristic X-radiation. As a result, each of thesesamples was found an amorphous alloy. The material alloy amorphizedeasily because the alloy contained B at a relatively high concentration.

[0166] Next, the rapidly solidified alloy samples Nos. 13 to 19 wereheat-treated within Ar gas. Specifically, the rapidly solidified alloyswere kept at the respective heat treatment temperatures shown on therightmost column in Table 3 for 6 minutes and then cooled to roomtemperature. Thereafter, the magnetic properties of these samples weremeasured using a vibrating sample magnetometer. The results are shown inthe following Table 4: TABLE 4 Magnetic properties H_(cJ) (BH)_(max)B_(r) (T) (kA/m) (kJ/m³) 13 0.83 957 111 14 0.79 906 105 15 0.82 826 10416 0.70 1073 78 17 0.63 197 28 18 0.71 462 56 19 1.0 30 12

[0167] As can be seen from Table 4, samples Nos. 13 to 16 representingexamples of the present invention showed magnetic properties that weresuperior to those of samples Nos. 17 to 19 representing the referenceexamples.

[0168]FIG. 9 illustrates the demagnetization curves of samples Nos. 13and 17. In FIG. 9, the ordinate represents the remanence while theabscissa represents the demagnetization field strength. As can be seenfrom FIG. 9, the demagnetization curve of sample No. 13 has loopsquareness much better than that of sample No. 17. FIG. 10 illustratesthe XRD patterns of sample No. 13 before and after the heat treatment,while FIG. 11 illustrates the XRD patterns of sample No. 17 before andafter the heat treatment.

[0169] As can be seen from FIG. 10, where Ti was added, no diffractionpeaks representing crystallinity were observed in the alloy yet to beheat-treated (i.e., in the “as-spun” state). However, after the alloyhad been heat-treated at 660° C. for 6 minutes, some diffraction peaks,showing the existence of a compound phase with the Nd₂Fe₁₄B crystalstructure, were observed. In this case, diffraction peaks correspondingto the α-Fe phase were also observed but the intensities thereof werenot so high. And when the heat treatment was conducted at 780° C., thediffraction peaks corresponding to the α-Fe phase increased theirintensities. Accordingly, the crystallization temperature of the α-Fephase would be higher than that of the Nd₂Fe₁₄B phase.

[0170] On the other hand, where no Ti was added, no diffraction peaks,showing the existence of a compound phase with the Nd₂Fe₁₄B crystalstructure, were observed, but a diffraction peak corresponding to theα-Fe phase was clearly identified after the alloy had been heat-treatedat 660° C. for 6 minutes. This can be easily seen from the results shownin FIG. 11. Also, this result shows that the α-Fe phase had nucleatedand grown earlier than the crystallization of the Nd₂Fe₁₄B phase. Andwhere the heat treatment was conducted at 780° C., the diffraction peakof the α-Fe phase increased its intensity considerably. Thus, it can beseen that the grain size of the α-Fe phase increased excessively.

[0171] As can be seen from these results, where the mole fraction x of Qis 15 at % or more, the mole fraction z of M is preferably greater than3.0.

[0172] Next, a melt of an alloy with a compositionNd₉Fe_(78.7)B_(10.3)Ti₂ (where the mole fractions are indicated inatomic percentages) was rapidly cooled with the pressure of the ambientgas and the surface velocity of the roller changed.

[0173] The quartz crucible used for preparing the melt had an orificewith a diameter of 0.8 mm at the bottom. Accordingly, the alloyincluding these materials was melted in the quartz crucible to be amelt, which was then dripped down through the orifice. The materialalloy was melted by a high frequency heating method within an argonenvironment at a pressure of 1.33 kPa. In the illustrated examples, thetemperature of the melt was set to 1500° C.

[0174] The surface of the melt was pressurized with Ar gas at 26.7 kPa,thereby propelling the melt against the outer circumference of a copperchill roller, which was located 0.7 mm under the orifice. The otherconditions were substantially the same as those placed on sample Nos. 1to 19. In these examples, the pressure of the quenching environment,roller surface velocity and heat treatment temperature were changed asshown in the following Table 5: TABLE 5 Roller Heat Environment surfacetreatment pressure velocity temperature (kPa) (m/s) (° C.) 20 40.0 10.0620 21 35.0 15.0 640 22 40.0 20.0 650 23 80.0 23.0 660 24 60.0 12.0 64025 40.0 28.0 690 26 10.0 15.0 680 27 40.0 35.0 700 28 40.0 5.0 600

[0175] The rapidly solidified alloy structures obtained by these rapidcooling processes were analyzed using Cu—Kα characteristic X-radiation.Using a TEM, the present inventors confirmed that each of samples Nos.20 to 25 contained 60 volume % or more of Nd₂Fe₁₄B phase and alsoconfirmed the existence of the α-Fe and Fe₂₃B₃ phases in addition to theNd₂Fe₁₄B phase. FIG. 12 illustrates XRD patterns of sample No. 21. InFIG. 12, the profile labeled as “as-spun” is the XRD pattern of arapidly solidified alloy yet to be heated. FIG. 12 also illustrates anXRD pattern of sample No. 21 that had already been heated (see thefollowing description).

[0176] As for sample No. 26, diffraction peaks, showing the existence ofthe Nd₂Fe₁₄B, α-Fe and Fe₂₃B₆ phases, were confirmed. However, only halopatterns were observed for sample No. 27. And for sample No. 28, anintense diffraction peak corresponding to the α-Fe phase and barelyrecognizable diffraction peaks of the Nd₂Fe₁₄B phase were observed. Itshould be noted that a lot of amorphous phase existed in sample No. 26.

[0177] Next, the rapidly solidified alloy samples Nos. 20 to 26 wereheat-treated within Ar gas. Specifically, the rapidly solidified alloyswere kept at the respective heat treatment temperatures shown on therightmost column in Table 5 for 6 minutes and then cooled to roomtemperature. Thereafter, the magnetic properties of these samples weremeasured using a vibrating sample magnetometer. The results are shown inthe following Table 6: TABLE 6 Magnetic properties B_(r) (T) H_(cj)(kA/m) (BH)_(max) (kJ/m³) 20 0.89 705 124 21 0.94 650 130 22 0.95 600126 23 0.88 683 124 24 0.90 670 125 25 0.87 588 120 26 0.83 780 114 270.81 754 103 28 0.64 334 38

[0178] As can be seen from Table 6, each of samples Nos. 20 to 25exhibited excellent hard magnetic properties including a remanence B_(r)of 0.85 T or more, an intrinsic coercivity H_(cJ) of 480 kA/m or moreand a maximum energy product (BH)_(max) of 120 kJ/m³ or more.

[0179]FIG. 13 illustrates the demagnetization curves of samples Nos. 21and 26, respectively. In FIG. 13, the ordinate represents the remanencewhile the abscissa represents the demagnetization field strength. As canbe seen from FIG. 13, the demagnetization curve of sample No. 21 hasloop squareness of the demagnetization curves much better than that ofsample No. 26. The sample No. 26 showed deteriorated loop squarenessprobably because the crystal grain size was large

[0180] Next, each of the heat-treated samples Nos. 20 to 25 was analyzedusing Cu—Kα characteristic X-radiation to identify the respectiveconstituent phases thereof. The size of the constituent phases weredetermined by TEM. As a result, the R₂Fe₁₄B phase had an average crystalgrain size ranging from 20 nm to 100 nm and the α-Fe and iron-basedboride phases had an average crystal grain size ranging from 1 nm to 50nm.

[0181] As for samples Nos. 26 and 28 on the other hand, the types of theconstituent phases included did not change before and after the heattreatment. But for sample No. 27, the nucleation and growth of α-Fe andFe₂₃B₆ phases, as well as the R₂Fe₁₄B phase, was confirmed.

[0182] As can be seen from these results, the quenching environmentshould have a pressure of 30 kPa or more. And in that case, the surfacevelocity of the chill roller is preferably from 10 m/sec to 30 m/sec.

[0183] According to the present invention, a melt of a material alloy,containing Ti as an additive, is rapidly cooled and solidified, therebyreducing the amount of a rare earth element needed for a permanentmagnet. And yet the resultant magnet can exhibit excellent magneticproperties, or has sufficiently high coercivity and remanence.

[0184] Also, according to the present invention, even if a rapidlysolidified alloy is prepared by a rapid cooling process at a decreasedcooling rate, the addition of Ti can suppress the precipitation of theα-Fe phase during the rapid cooling process. Therefore, a strip castingmethod, or a rapid cooling process resulting in a relatively low coolingrate and suitably applicable to mass production, can be adopted in thepresent invention, thus reducing the production cost advantageously.

[0185] It should be understood that the foregoing description is onlyillustrative of the invention. Various alternatives and modificationscan be devised by those skilled in the art without departing from theinvention. Accordingly, the present invention is intended to embrace allsuch alternatives, modifications and variances which fall within thescope of the appended claims.

What is claimed is:
 1. An iron-based rare earth alloy magnet with acomposition represented by the general formula:(Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z), where T is at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare earth element substantially excluding La and Ce; and M is atleast one metal element selected from the group consisting of Ti, Zr andHf and always includes Ti, wherein the mole fractions x, y, z and m meetthe inequalities of: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1 at %≦z≦12at %; and 0≦m≦0.5, respectively, and wherein said magnet comprises twoor more ferromagnetic crystalline phases including hard and softmagnetic phases; an average grain size of the hard magnetic phase isbetween 10 nm and 200 nm; and an average grain size of the soft magneticphase is between 1 nm and 100 nm.
 2. The magnet of claim 1, wherein themole fractions x, y and z meet the inequalities of: 10 at %<x<17 at %; 8at %≦y≦9.3 at %; and 0.5 at %≦z≦6 at %, respectively.
 3. The magnet ofclaim 1, wherein R₂Fe₁₄B phase, boride phase and α-Fe phase coexist inthe same metal structure.
 4. The magnet of claim 3, wherein an averagecrystal grain size of the α-Fe and boride phases is between 1 nm and 50nm.
 5. The magnet of claim 4, wherein said boride phase comprises aniron-based boride with ferromagnetic properties.
 6. The magnet of claim5, wherein said iron-based boride comprises at least one of Fe₃B andFe₂₃B₆.
 7. The magnet of claim 1, wherein said mole fractions x and zmeet z/x≧0.1.
 8. The magnet of claim 1, wherein said mole fraction y ofthe rare earth element(s) R is 9.5 at % or less.
 9. The magnet of claim1, wherein said mole fraction y of the rare earth element(s) R is 9.0 at% or less.
 10. The magnet of one of claims 1 to 6, wherein said magnetsis shaped in a thin strip with a thickness of between 10 μm and 300 μm.11. The magnet of one of claims 1 to 6, wherein said magnets has beenpulverized into powder particles.
 12. The magnet of claim 11, wherein amean particle size of the powder particles is between 30 μm and 250 μm.13. The magnet of one of claims 1 to 6, wherein said magnet exhibits acoercivity H_(cJ) of 480 kA/m or more and a remanence B_(r) of 0.7 T ormore.
 14. The magnet of one of claims 1 to 6, wherein said magnetexhibits a remanence B_(r) of 0.85 T or more, a maximum energy product(BH)_(max) of 120 kJ/m³ or more and an intrinsic coercivity H_(cJ) of480 kA/m or more.
 15. A bonded magnet formed by molding a magnet powder,comprising the powder particles of the iron-based rare earth alloymagnet as recited in claim 11 and a resin binder.
 16. A rapidlysolidified alloy for an iron-based rare earth alloy magnet, the alloyhaving a composition represented by the general formula:(Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z), where T is at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare earth element substantially excluding La and Ce; and M is atleast one metal element selected from the group consisting of Ti, Zr andHf and always includes Ti, wherein the mole fractions x, y, z and m meetthe inequalities of: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1 at %≦z≦12at %; and 0≦m≦0.5, respectively.
 17. The alloy of claim 16, wherein saidalloy comprises a structure, in which substantially no α-Fe phase isincluded but R₂Fe₁₄B compound and amorphous phases are included and inwhich the R₂Fe₁₄B phase accounts for a 60 volume percent or more. 18.The alloy of claim 17, wherein said mole fractions x, y and z meet theinequalities of: 10 at %<x<17 at %; 8 at %≦y≦9.3 at %; and 0.5 at %≦z≦6at %, respectively, and wherein said R₂Fe₁₄B phase, accounting for 60volume percent or more of the alloy, has an average grain size of 50 nmor less.
 19. A rapidly solidified alloy for an iron-based rare earthalloy magnet, wherein the alloy is prepared by rapidly cooling a melt ofa material alloy comprising Fe, Q, R and Ti, where Q is at least oneelement selected from the group consisting of B and C; and R is a rareearth element, wherein the rapidly solidified alloy has a structure inwhich an amorphous phase is included and in which heat treatment startsto grow a compound crystalline phase with an R₂Fe₁₄B crystallinestructure before starting to grow an α-Fe crystalline phase.
 20. Amethod for producing an iron-based rare earth alloy magnet, the methodcomprising the steps of: preparing a melt of a material alloy thatincludes Fe, Q, R and Ti, where Q is at least one element selected fromthe group consisting of B and C, and R is a rare earth element; coolingthe melt to make a solidified alloy including an amorphous phase; andheating the solidified alloy to start growing a compound crystallinephase with an R₂Fe₁₄B crystalline structure and then an α-Fe crystallinephase.
 21. The method of claim 20, wherein said cooling step is effectedvia a strip casting process.
 22. A method for producing an iron-basedrare earth alloy magnet, comprising the steps of: preparing a melt of amaterial alloy, the material alloy having a composition represented bythe general formula: (Fe_(1-m)T_(m))_(100-x-y-z)Q_(x)R_(y)M_(z), where Tis at least one element selected from the group consisting of Co and Ni;Q is at least one element selected from the group consisting of B and C;R is at least one rare earth element substantially excluding La and Ce;and M is at least one metal element selected from the group consistingof Ti, Zr and Hf and always includes Ti, the mole fractions x, y, z andm meeting the inequalities of: 10 at %<x≦20 at %; 6 at %≦y<10 at %; 0.1at %≦z≦12 at %; and 0≦m≦0.5, respectively; rapidly cooling the melt tomake a rapidly solidified alloy comprising an R₂Fe₁₄B crystalline phase;and thermally treating the rapidly solidified alloy to form a structurein which two or more ferromagnetic crystalline phases, including hardand soft magnetic phases, exist, an average grain size of the hardmagnetic phase is equal to or greater than 10 nm and equal to or lessthan 200 μm, and an average grain size of the soft magnetic phase isequal to or greater than 1 nm and equal to or less than 100 nm.
 23. Themethod of claim 22, wherein said rapidly solidified alloy made in thecooling step includes an R₂Fe₁₄B phase at 60 volume percent or more. 24.The method of claim 22, wherein said cooling step comprises rapidlycooling the melt within an ambient gas at a pressure of 30 kPa or moreto make a rapidly solidified alloy including an R₂Fe₁₄B phase with anaverage grain size of 50 nm or less.
 25. The method of claim 24, whereintsaid cooling step comprises: bringing the melt into contact with thesurface of a rotating chill roller to obtain a supercooled liquid alloy;and dissipating heat from the supercooled alloy into the ambient gas togrow the R₂Fe₁₄B phase after the supercooled alloy has left the chillroller.
 26. The method of claim 22, comprising the step of forming astructure in which three or more crystalline phases, including at leastR₂Fe₁₄B compound, α-Fe and boride phases, are included, an averagecrystal grain size of the R₂Fe₁₄B phase is between 20 nm and 150 nm, andan average crystal grain size of the α-Fe and boride phases is between 1nm and 50 nm.
 27. The method of claim 26, wherein said steps of heatingand crystallizing the rapidly solidified alloy produces a boride phasecomprising an iron-based boride with ferromagnetic properties.
 28. Themethod of claim 27, wherein said steps of heating and crystallizing therapidly solidified alloy produces a boride phase comprising aniron-based boride comprises Fe₃B and/or Fe₂₃B₆.
 29. The method of claim22, wherein said step of cooling the melt is effected via a stripcasting process.
 30. A method for producing a bonded magnet, the methodcomprising the steps of: preparing a powder of the iron-based rare earthalloy magnet by the method as recited in one of claims 22 to 29; andprocessing the powder into a bonded magnet.
 31. The magnet of claim 11,wherein the powder particles have surfaces that have been subjected to asurface treatment.
 32. The bonded magnet of claim 15, comprising asurface that has been subjected to a surface treatment.